Rare earth magnet alloy ingot, manufacturing method for the same, R-T-B type magnet alloy ingot, R-T-B type magnet, R-T-B type bonded magnet, R-T-B type exchange spring magnet alloy ingot, R-T-B type exchange spring magnet, and R-T-B type exchange spring bonded magnet

ABSTRACT

One object of the present invention is to provide a rare earth magnet alloy ingot, which has improved magnetic properties. In order to achieve the object, the present invention provides a rare earth magnet alloy ingot, wherein the rare earth magnet alloy ingot comprises an R-T-B type magnet alloy (R represents at least one element selected from among rare earth elements, including Y; and T represents a substance predominantly comprising Fe, with a portion of Fe atoms being optionally substituted by Co, Ni, Cu, Al, Ga, Cr, and Mn) containing at least one element selected from among Nd, Pr, and Dy in a total amount of 11.8 to 16.5% by atom and B in an amount of 5.6 to 9.1% by atom; and wherein as determined in an as-cast state of the alloy ingot, R-rich phase that measures 100 μm or more is substantially absent on a cross section.

CROSS REFERENCE TO RELATED APPLICATIONS

This is a divisional of application Ser. No. 10/232,520 filed Sep. 3,2002, now U.S. Pat. No. 7,014,718, which claims benefit of U.S.Provisional Application No. 60/322,748 filed Sep. 18, 2001, U.S.Provisional Application No. 60/367,720 filed Mar. 28, 2002 and U.S.Provisional Application No. 60/396,754 filed Jul. 19, 2002.

BACKGROUND OF THE INVENTION

1. Field of the Invention

The present invention relates to a rare earth alloy ingot, a sinteredmagnet comprising the rare earth alloy ingot, a production method for arare earth alloy ingot, a production method for a rare earth alloyflake, an R-T-B type magnet alloy ingot, an R-T-B type magnet, an R-T-Btype magnet alloy flake, an R-T-B type magnet alloy powder, an R-T-Btype bonded magnet, an R-T-B type exchange spring magnet alloy ingot, anR-T-B type exchange spring magnet, an R-T-B type exchange spring magnetalloy powder, and an R-T-B type exchange spring bonded magnet.

2. Description of the Related Art

In recent years, production of Nd—Fe—B alloys serving as magnet alloyshas sharply increased by virtue of high-performance characteristics ofthe alloys, and these alloys are employed in HDs (hard disks), MRI(magnetic resonance imaging), a variety of motors, etc. Typically, aportion of Nd atoms is substituted by another rare earth element such asPr or Dy (as used herein, Nd and the substituted Nd are referred to asR, moreover, Y is at least one selected from the rare earth elementscontaining Y) or a portion of Fe is substituted by another metal elementsuch as Co, Ni, Cu, Ga or Al (as used herein, Fe and the substituted Feare referred to as T). Such substituted alloys as well as Nd—Fe—B alloysare generally referred to as R-T-B type alloys.

An R-T-B alloy contains, as the dominant phase, a ferromagnetic phaseformed of R₂T₁₄B crystals, which contribute to magnetization, and, ingrain boundaries of the R₂T₁₄B crystals, a nonmagnetic R-rich phasehaving a low melting point and containing a rare earth element(s) athigh concentration. The R-T-B alloy is an active metallic material.Therefore, the alloy is generally melted and mold-cast in vacuum orunder inert gas.

In a typical method of producing a magnet, an ingot of the alloy ispulverized to powder having a particle size of about 3 μm (as measuredby means of FSSS (Fisher Sub-Sieve Sizer)); the powder is subjected topress-forming in a magnetic field; the resultant compact is sintered ina sintering furnace at a temperature as high as about 1,000 to about1,100° C.; and in accordance with needs, the sintered product is heated,mechanically processed, and plated for corrosion prevention.

The R-rich phase plays the following important roles.

-   (1) Since the R-rich phase has a low melting point, the phase    liquefies during sintering, thereby contributing to achievement of    high density of the resultant magnet, leading to improved    magnetization-   (2) The R-rich phase functions to smoothen grain boundaries, thereby    reducing the number of nucleation sites in a reversed magnetic    domain, thereby enhancing the coercivity.-   (3) The R-rich phase magnetically insulates the dominant phase,    thereby enhancing the coercivity.

Thus, attainment of a uniformly dispersed R-rich phase is critical,because otherwise magnet characteristics of the produced magnet areadversely affected

The distribution of the R-rich phase in a magnet—the finalproduct—depends greatly on the metallographic structure of the rawmaterial alloy ingot. Specifically, when the alloy is mold-cast, a slowcooling rate often results in formation of large crystal grains. In sucha case, the particle size of the pulverized product becomes considerablysmaller than that of the crystal grain size. When the alloy ismold-cast, R-rich phase is included not in crystal grains but virtuallyin crystal grain boundaries. Therefore, particles formed only of thedominant phase containing no R-rich phase and those formed only of theR-rich phase result, making it difficult to mix the dominant phase andR-rich phase homogeneously.

Another problem involved in mold cat is that γ-Fe tends to be formed asprimary crystals, due to the slow cooling rate. At approximately 910° C.or lower, γ-Fe transforms into α-Fe, which deteriorates pulverizationefficiency dung production of magnets. If α-Fe remains even aftersintering, magnetic characteristics of the sintered product aredeteriorated. Thus, the ingot obtained through mold casing must besubjected to homogenization treatment at high temperature for a longperiod of time in order to remove α-Fe.

In order to solve the above problems, the strip casting methodhereinafter referred to as the SC method), which ensures a cooling rateduring casing faster than that attainable by mold casting, is proposedand employed in actual production steps.

In the SC method, a molten metal is slowly poured onto a copper rollwhose inside is cooled by water and which rotates at a peripheralvelocity of about 1 mm/sec, and is solidified through rapid cooling, tothereby produce flake having a thickness in a range from about 0.1 to 1mm (Japanese Patent Application Laid-Open (kokai) Nos. 05-222488 and05-295490). During casting, the molten metal was solidified throughrapid cooling, to thereby yield an alloy having a microcrystallinestructure in which R-rich phase is minutely dispersed. Since the R-richphase is minutely dispersed in the alloy, dispersion of R-rich phase inthe product obtained by pulverizing and sintering the alloy becomes alsosatisfactory, to thereby successfully attain improved magneticcharacteristics (Japanese Patent Application Laid-Open (kokai) Nos.5-222488 and 5-295490). However, even when the above method is employed,α-Fe is unavoidably formed as the R content (%) decreases. For example,when the Nd content of an Nd—Fe—B ternary alloy is 28% by weight orless, α-Fe generation becomes significant

The thus-formed α-Fe considerably deteriorates pulverizability of analloy ingot in magnet production steps.

FIG. 9 is a back-scattered electron image, observed under an SEM(scanning electron microscope), showing a cross section of an Nd—Fe—Bingot (Nd: 30.0% by mass) cast through a conventional SC method.

In FIG. 9, Nd-rich (i.e., R of the R-rich phase is Nd) phase correspondsto bright portions. Some portion of the Nd-rich phase assumes the shapeof linked rods extending in the solidification direction (left (rollside) to right (free side)). Another portion of the Nd-rich phaseassumes a dot shape and is dispersed In the rod-shaped Nd-rich phase,the growth direction in grain bodies and that in crystal grains coincidewith the longitudinal direction of the rod-shaped Nd-rich phase.Although the rod-shaped phase is slightly reduced or fragmented throughheat treatment performed after casting, effects exerted during castingstill prevail, and the dot-shaped or rod-shaped Nd-rich phase showsnonuniform dispersion. Such a microcrystalline feature is typical to across-sectional metallographic structure of an Nd—Fe—B alloy ingot castthrough the SC method.

As explained above, the R-T-B type alloy contains, as the dominantphase, a ferromagnetic phase formed of R₂T₁₄B crystals, which contributeto magnetization, and, in grain boundaries of the R₂T₁₄B crystals, anonmagnetic R-rich phase having a low melting point and containing arare earth element(s) at high concentration. The R-T-B type alloy is anactive metallic material. Therefore, the alloy is generally melted andcast in vacuum or under inert gas, and the cast alloy provides sinteredmagnets and bonded magnets. Below, the sintered magnet and the bondedmagnets are explained.

(1) Sintered Magnet

Alloy ingots for sintered magnets are produced through, among othermethods, the book molding method (hereinafter referred to as the BMmethod) and the SC method. In the BM method, a molten metal is cast in acopper mold or an iron mold whose inside is cooled by water, to therebyproduce an ingot having a thickness of about 5 to about 50

The alloy ingot produced through any of the above methods is pulverizedin an inert gas atmosphere, such as argon, nitrogen, to have a particlesize of about 3 μm (as measured by means of an FSSS (Fisher Sub-SieveSizer)); the resultant powder is subjected to press-forming in amagnetic field at 0.8 to 2 ton/cm²; the resultant compact is sintered ina sintering furnace at a temperature as high as about 1,000 to about1,100° C. (hereinafter, the steps of pulverization to sintering arecollectively referred to as the powder metallurgical method); and inaccordance with needs, the sintered product is heated at 500 to 800° C.,mechanically processed, and plated for prevention of corrosion, tothereby produce a magnet.

Among these methods, the SC method provides a minute microcrystallinestructure and forms an alloy in which a low-melting-temperature R-richphase formed of concentrated nonmagnetic rare earth elements is minutelydispersed. Since the R-rich phase is minutely dispersed in the alloy,dispersibility of the R-rich phase after pulverizing and sintering thealloy also becomes satisfactory, to thereby successfully attain improvedmagnetic characteristics as compared with those of alloy ingots producedthrough the BM method.

(2) Bonded Magnet

An alloy ingot for bonded magnets, in the form of ribbon having athickness in a range from 10 to 100 μm, is produced through theultra-rapid-cooling method; i.e., by injecting a molten metal from acrucible, via an orifice provided in the bottom of the crucible, onto acopper roll which rotates at a high peripheral velocity of about 20m/sec. The ribbon produced through the ultra-rapid-cooling method may beheated at 400 to 1,000° C. in accordance with needs, followed bypulverization to powder having a particle size of 500 μm or less. Amixture of the powder and a resin is press-molded or injection-molded,to thereby form a magnet. Since the ribbon is isotropic in terms ofmagnetic characteristics, the bonded magnet produced from the ribbonalso exhibits magnetic isotropy.

Recently, there has been proposed an exchange spring magnet having acomposite structure of a hard magnetic phase and a soft magnetic phase,each phase comprising crystal grains in a range from 10 to 100 nm insize. An alloy ingot for exchange spring magnets, containingconsiderably minute crystal grains, is generally produced through theultra-rapid-cooling method. The produced ingot may be heated at 400 to1,000° C. in accordance with needs, followed by pulverization to powderhaving a particle size of 500 μm or less. A mixture of the powder and aresin is press-molded or injection-molded, to thereby form an exchangespring magnet. In the exchange spring magnet, residual magnetic fluxdensity and coercive force are generally determined by crystal grains ofthe soft magnetic phase and crystal grains of the hard magnetic phase,respectively. Since the hard magnetic phase of the exchange springmagnet must exhibit a highly anisotropic magnetic field, the hardmagnetic phase is formed of a rare earth material such as R₂T₁₄B,Sm₁Co₅, or Sm₂Co₁₇. The soft magnetic phase is formed of Fe, Fe₂B, Fe₃B,etc., which exhibit high saturation magnetization.

In an as-cast state, R-T-B type magnet alloy ingots produced through theBM method or the SC method exhibit very weak magnetic characteristics,and so, they cannot be used as a magnet. The reason therefor is asfollows. In the case of R-T-B type magnets, coercive force is exhibitedon a nucleation-based mechanism. Specifically, crystal grin boundariescontain lattice defects and irregularities in an as-cast state, andthese lattice defects and irregularities serve as nuclei for generatinga reverse magnetic domain (hereinafter the nuclei are referred to asnucleation sites). Even when a weak reverse magnetic field is applied,magnetization inversion occurs from the nucleation sites, resulting inmagnetization inversion of the entirety of crystal gas. In particular,an alloy ingot produced through the BM method contains a large number ofcrystal grams having a major grain size of about some mm, and an alloyingot produced through the SC method contains a large number of crystalgrains having a major grain size of 100 μm or more. By virtue of havingsuch a large grain size, the volume required for magnetization inversionwith respect to the total volume of the alloy is large, resulting invery poor magnetic characteristics.

To avoid this, as described above, the alloy ingot is pulverize to havea particle size of about 3 Sun, followed by sintering, to therebyproduce a magnet. The thus-produced magnet has a crystal grain size ofabout 5 to about 20 μm, and the low-melting-temperature R-rich phasewhich becomes a liquid phase during sintering smoothens irregularitiesof grain boundaries, leading to reduction of nucleation sites, therebyenhancing coercive force. However, the steps of pulverization tosintering involve a considerably high cost Particularly when the alloypowder is an active R-T-B type magnet alloy powder, measures such asperforming the steps of pulverization to sintering in an inert gasatmosphere are required from the viewpoints of product quality with lessvariation and greater safety in production steps. Such measures alsoincrease the cost.

Meanwhile, ribbon for R-T-B type bonded magnets produced through theultra-rapid-cooling method is heated at 500 to 800° C. in accordancewith needs, so as to obtain optical magnetic characteristics. Throughheat treatment, ribbon having a crystal grain size in a range from 10 to100 nm and exhibiting magnetic isotropy is provided Since theribbon-form ingot is not practical for use, the ribbon is pulverized tohave a particle size of 500 μm or less. A mixture of the powder and aresin is press-molded or injection-molded, to thereby provide anisotropic bonded magnet There has also been proposed a method forproducing bulk isotropic magnets including hot-pressing the ribbon at700° C. and 1 ton/cm² (R. W. Lee, Appl. Phys. Lett. 46 (1985), JapanesePatent Application laid Open (kokai) No. 60-100402).

However, as compared with the BM method and the SC method, theultra-rapid-cooling method has low productivity. In addition, aproduction method of bulk isotropic magnets including hot-pressingrequires a high cost.

The alloy ribbon for exchange spring magnets produced through theultra-rapid-cooling method is also heated at 500 to 800° C. inaccordance with needs, so as to obtain optimal magnetic characteristics.Through heat treatment, ribbon having a crystal grain size in a rangefrom 10 to 100 nm and exhibiting magnetic isotropy is provided. Sincethe ribbon-form ingot is not practical for use, the ribbon is pulverizedto have a particle size of 500 μm or less. A mixture of the powder and aresin is press-molded or injection-molded, to thereby provide anisotropic bonded magnet. There has also been disclosed a method forproducing bulk isotropic magnets including plasma sintering the ribbon(SPS method) (e.g., Ono, Waki, Fujiki, Shimada, Yamamoto, Sonoda, &Tani, Resume of Lectures, Convention of The Japan Institute of Metals,spring, 2000).

However, as described above, productivity of the ribbon throughultra-rapid-cooling method is low. In addition, a production method ofbulk isotropic magnets including plasma sintering involves asignificantly high cost.

The present inventors previously improved conventional centrifugalcasting methods and devised another solidification process and anapparatus therefor (Japanese Patent Application Laid-Open (kokai) Nos.08-13078 and 08-332557). Specifically, molten metal is introduced into arotating mold via a box-like tundish, which is disposed in areciprocative manner inside the mold and has a plurality of nozzles,whereby the molten metal is deposited and solidified on the innersurface of the rotating mold (this process is called a CC (CentrifugalCasting) process).

In the CC process, molten metal is continuously poured onto an ingotwhich has already been deposited and solidified. The additionally castmolten metal semi-solidifies while the mold makes one rotation, wherebythe rate of solidification can be increased. However, in the productionof an alloy of low R content through the CC method, α-Fe which isdetrimental to magnetic characteristics and magnet production steps isunavoidably formed due to low cooling rate in a high-temperature zone.

In order to prevent formation of α-Fe in R-T-B type alloy ingots, thepresent inventors attempted to increase the solidification-cooling ratein the CC process by reducing the deposition rate of a molten metal andpreviously proposed a centrifugal casting method including sprinkling amolten metal from a rotating tundish and causing the sprinkled moltenmetal to be deposited on an inner surface of a rotating mold (JapanesePatent Application No. 2000-262605). Through employment of the abovemethod, formation of α-Fe was found to be suppressed Thereby, a castalloy ingot of low R content, which enhances magnetic characteristics ofproduced magnets, can be produced.

However, when the R content decreases, R-rich phase content decreases,possibly resulting in failure to produce sintered magnets of highdensity and enhanced coercive force. Therefore, it is thought that aminute and uniform dispersion state of the R-rich phase must be attainedthrough more rapid cooling-solidification so as to attain furtherenhanced magnetic characteristics.

In addition, the thus-produced R-T-B type alloy ingot contains a largenumber of crystal grains having a major size of 1,000 μm or more, andexhibits very poor magnetic characteristics in an as-cast state.Therefore, further enhancement of the solidification-cooling rate, tothereby reduce the crystal grain size, is deemed necessary.

The present inventors have carried out extensive studies on improvementof conventional centrifugal casting methods and have invented a methodwhich controls the rate of feeding a molten metal and raises the heattransfer efficiency from the cast surface of the alloy ingot which hasbeen deposited and solidified to the inner surface wall of the castingmold.

Thereby, it was confirmed that alloy ingots in which R-rich phase isminutely and uniformly dispersed and which have not been conventionallyproduced can be obtained, and that sintered magnets produced from theingots exhibit excellent magnetic characteristics.

In addition, thereby, it is possible to obtain a cast ingot of an R-T-Btype alloy having fine crystal grains not available in the past, and ithas been confirmed that the cast ingot, as it is, exhibits excellentisotropic magnetic properties.

The present invention has an object of providing a production method fora rare earth alloy ingot and a production method for a rare earth alloyflake which improve the efficiency of heat transfer from the castsurface of the cast ingot to the inner surface wall of the casting mold.

In addition, the present invention has an object of providing a rareearth magnet alloy ingot and a sintered magnet, which have improvedmagnetic properties.

In addition, the present invention has an object of providing an R-T-Btype magnet alloy ingot, an R-T-B type magnet, an R-T-B type magnetalloy flake, an R-T-B type magnet alloy powder, an R-T-B type bondedmagnet, an R-T-B type exchange spring magnet alloy ingot, an R-T-B typeexchange spring magnet, an R-T-B type exchange magnet alloy powder, andan R-T-B type exchange spring bonded magnet, which have fine crystalgrains not available conventionally.

SUMMARY OF THE INVENTION

In order to achieve the object, the present invention provides a rareearth magnet alloy ingot, characterised by comprising an R-T-B typemagnet alloy (R represents at least one element selected from among rareearth elements, including Y; and T represents a substance predominantlycomprising Fe, with a portion of Fe atoms being optionally substitutedby Co, Ni, Cu, Al, Ga, Cr, and Mn.) containing at least one elementselected from among Nd, Pr, and Dy in a total amount of 11.8 to 16.5% byatom and B in an amount of 5.6 to 9.1% by atom, and characterized inthat, as determined in an as-cast state of the alloy ingot, R-rich phasethat measures 100 μm or more is substantially absent on a cross section.

In addition, in order to achieve the object, the present inventionprovides another rare earth magnet alloy ingot, characterized bycomprising an R-T-B type magnet alloy (R represents at least one elementselected from among r earth elements, including Y; and T represents asubstance predominantly comprising Fe, with a portion of Fe atoms beingoptionally substituted by Co, Ni, Cu, Al, Ga, Cr, and Mn.) containing atleast one element selected from among Nd, Pr, and Dy in a total amountof 11.8 to 16.5% by atom and B in an amount of 5.6 to 9.1% by atom, andcharacterized in that, as determined in an as-cast state of the alloyingot, an area in which R-rich phase that measures 50 μm or less isdispersed accounts for at least 50% the cross section.

In addition, in order to achieve the object, the present inventionprovides another rare earth magnet alloy ingot, characterized bycomprising an R-T-B type magnet alloy (R represents at least one elementselected from among rare earth elements, including Y; and T represents asubstance predominantly comprising Fe, with a portion of Fe atoms beingoptionally substituted by Co, Ni, Cu, Al, Ga, Cr, and Mn.) containing atleast one element selected from among Nd, Pr, and Dy in a total amountof 11.8 to 16.5% by atom and B in an amount of 5.6 to 9.1% by atom, andcharacterized in that, as determined in an as-cast state of the alloyingot, R-rich phase having an aspect ratio of at least 20 issubstantially absent on a cross section.

In the rare earth magnet alloy ingots, it is preferable that the crystalgrains having a diameter of at least 1,000 μm as measured along themajor axis occupy an area percentage of at least 5%, and average R-richphase spacing is 10 μm or less.

In the rare earth magnet alloy ingots, it is also preferable that α-Feis substantially absent.

In the rare earth magnet alloy ingots, it is preferable that the rareearth magnet alloy ingot is cast through centrifugal casting comprisingreceiving molten metal by means of a rotary body, sprinkling the moltenmetal by the effect of rotation of the rotary body; and causing thesprinkled molten metal to be deposited and solidify on an inner surfaceof a rotating cylindrical mold, the inner surface including a non-smoothsurface.

In the rare earth magnet alloy ingot, it is preferable that an axis ofrotation of the rotary body and an axis of rotation of the cylindricalmold form an angle of inclination θ.

In addition, in order to achieve the object, the present inventionprovides a sintered magnet produced from the rare earth magnet alloyingot as a raw material.

In addition, in order to achieve the object, the present inventionprovides a method for producing a rare earth magnet alloy ingot,characterized by comprising receiving molten metal by means of a rotarybody; sprinkling the molten metal by the effect of rotation of therotary body; and causing the sprinkled molten metal to be deposited andsolidify on an inner surface of a rotating cylindrical mold, the innersurface including a non-smooth surface.

In the method for producing a rare earth magnet alloy ingot, it ispreferable that an axis of rotation of the rotary body and an axis ofrotation of the cylindrical mold form an angle of inclination θ.

In the method for producing a rare earth magnet alloy ingot, it is alsopreferable that the rare earth magnet alloy ingot is an R-T-B typemagnet alloy ingot.

In addition, in order to achieve the object, the present inventionprovides another method for producing a rare earth magnet alloy ingotcomprising receiving a molten alloy of rare earth metal alloy by meansof a rotary body, sprinkling the molten alloy by the effect of rotationof the rotary body, and causing the sprinkled molten alloy to bedeposited and solidify on an inner wall surface of a rotatingcylindrical mold; wherein a film having a thermal conductivity lowerthan that of material comprising the mold is provided to the inner wallsurface of the cylindrical mold.

In the production method, it is preferable for the inner wall surface ofthe rotating cylindrical mold to contain a non-smooth surface.

In the production method, it is preferable for the thermal conductivityof the film to equal 80 W/mK or less.

In the production method, it is preferable for the film to be made of ametal, a ceramic, or a metal-ceramic composite.

In the production method, it is preferable for the film to be providedon the inner wall surface of the mold by at least one selected fromcoating, plating, spray coating, and welding.

In the production method, it is preferable for the film to have athickness falling within a range of 1 μm to 1 mm.

In the production method, it is preferable for an axis of rotation ofthe rotary body and an axis of rotation of the cylindrical mold to forman angle of inclination θ.

In the production method, it is preferable that two or more layerscomprising rare earth alloy ingot are deposited and casted on the innerwall surface of the mold by centrifugal casting method.

In the production method, it is preferable to further comprise the stepof hot-working the obtained rare earth alloy ingot at 500 to 1,100° C.

In the production method, it is preferable to further comprise the stepof hot-treating the obtained rare earth alloy ingot at 400 to 1,000° C.

In the production method, it is preferable to further comprise the stepsof heat-treating the obtained rare earth alloy ingot at 1,000 to 1,100°C., and subsequently, an heat-treating at 400 to 1,000° C.

In the production method, it is preferable for the rare earth alloyingot to be an R-T-B type magnet alloy (R represents at least oneelement selected from among rare earth elements, including Y; and Trepresents a substance predominantly comprising Fe, with a portion of Featoms being optionally substituted by Co, Ni, Cu. A, Ga, Cr, and Mn.).

In addition, in order to achieve the object, the present inventionprovides a production method for a rare earth alloy flakes comprisingthe steps of: receiving a molten alloy of rare earth metal alloy bymeans of a rotary body; sprinkling the molten alloy by the effect ofrotation of the rotary body, and causing the sprinkled molten alloy tobe deposited and solidify on an inner wall surface of a rotatingcylindrical mold; wherein a film having a thermal conductivity lowerthan that of material comprising the old is provided to the inner wallsurface of the cylindrical mold; and casting is performed while alloyflakes deposited on the inner wall surface of the cylindrical mold arescraped.

In the production method, it is preferable to further comprise the stepof hot-treating the obtained rare earth alloy flakes at 400 to 1,000° C.

In the production method, it is preferable to further comprise the stepsof heat-treating the obtained rare earth alloy flakes at 1,000 to 1,100°C., and subsequently, an heat-treating at 400 to 1,000° C.

In the production method, it is preferable for the rare earth alloyflakes to be R-T-B type magnet alloy flakes (R represents at least oneelement selected from among rare earth elements, including Y; and Trepresents a substance predominantly comprising Fe, with a portion of Featoms being optionally substituted by Co, Ni, Cu, Al, Ga, Cr, and Mn.).

In addition, in order to achieve the object, the present inventionprovides an R-T-B type magnet alloy ingot, wherein it comprises at leastone element selected from among Nd, Pr, and Dy in a total amount of 11.8to 16.5% by atom and B in an amount of 5.6 to 9.1% by atom, with abalance being T (T represents a substance predominantly comprising Fe,with a portion of Fe atoms being optionally substituted by Co, Ni, Cu,Al, Ga, Cr, and Mn.); wherein it contains crystal grin having a grainsize of 10 μm or less in a volume of at least 50% on the basis of theentire volume of the alloy; and wherein it is produced by the productionmethod for an R-T-B type magnet alloy ingot.

In the R-T-D type magnet alloy ingot, it is preferable to have athickness of at least 1 mm as it is casted.

In addition, in order to achieve the object, the present inventionprovides an R-T-B type magnet, wherein it is obtainable by mechanicalprocessing at least one method selected from cutting, grinding,polishing, and blanking the R-T-B type magnet alloy ingot.

In the R-T-B type magnet, it is preferable to have a cylindrical shapehaving an outer diameter of at least 100 mm.

In addition, in order to achieve the object the present inventionprovides an R-T-B type magnet alloy flakes, wherein they arm produced bythe production method for the F-T-B type magnet alloy flakes; whereinthey comprise at least one element selected from among Nd, Pr, and Dy ina total amount of 11.8 to 16.5% by atom and B in an amount of 5.6 to9.1% by atom, with a balance being T (T represents a substancepredominantly comprising Fe, with a portion of Fe atoms being optionallysubstituted by Co, Ni, Cu, Al, Ga, Cr, and Mn.); and wherein theycontain crystal grains having a grain size of 10 μm or less in a volumeof at least 50% on the basis of the entire volume of the flakes.

In the R-T-B type magnet alloy flakes, it is preferable to have maximumlength of 5 cm or less, and the thickness of 1 mm or less.

In addition, in order to achieve the object, the present inventionprovides an R-T-B type magnet alloy powder produced throughpulverization of the R-T-B type magnet alloy ingot to a particle size of500 μm or less.

In addition, in order to achieve the object, the present inventionprovides an R-T-B type magnet alloy powder produced throughpulverization of the R-T-B type magnet alloy flakes to a particle sizeof 500 μm or less.

In addition, in order to achieve the object, the present inventionprovides an R-T-B type bonded magnet produced by use of the R-T-B typemagnet alloy powder.

In addition, in order to achieve the object, the present inventionprovides an R-T-B type exchange spring magnet alloy ingot, wherein it isproduced by the production method for rare earth alloy ingot; itcomprises at least one element selected from Nd, Pr, and Dy in a totalamount of 1 to 12% by atom and B in an amount of 3 to 30% by atom, witha balance being T (T represents a substance predominantly comprising Fe,with a portion of Fe atoms being optionally substituted by Co, Ni, Cu,Al, Ga, Cr, and Mn.); it is produced through formation of a composite ofcrystal grains of a hard magnetic phase and crystal grains of a softmagnetic phase; and it contains crystal grains of a hard magnetic phasehaving a grain size of 1 μm or less and crystal grains of a softmagnetic phase having a grain size of 1 μm or less in a volume of atleast 50% on the basis of the entire volume of the alloy.

In the R-T-B type exchange spring magnet alloy ingot, it is preferableto have a thickness of 1 mm or greater as it is casted.

In the R-T-B type exchange spring magnet alloy ingot, it is preferableto be heat treated at 400 to 1,000° C. after casting.

In addition, in order to achieve the object, the present inventionprovides an R-T-B type exchange spring magnet, wherein it is obtainableby mechanical processing at least one method selected from cutting,grinding, polishing, and blah, the R-T-B type exchange spring magnetalloy ingot.

In the R-T-B type exchange spring magnet, it is preferable to have acylindrical shape having an outer diameter of at least 100 mm. Inaddition, in order to achieve the object, the present invention providesan R-T-B type exchange spring magnet alloy powder produced throughpulverization of the R-T-B type exchange spring magnet alloy ingot to aparticle size of 500 μm or less.

Furthermore, in order to achieve the object, the present inventionprovides an R-T-B type exchange spring bonded magnet produced by use ofthe R-T-B type exchange spring magnet alloy powder.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a view showing a metallographic structure (cross section) of arare earth alloy ingot of the present invention.

FIG. 2 is a view showing an exemplary apparatus for producing a rareearth magnet alloy ingot according to the present invention.

FIG. 3 is a cross-sectional view of an exemplary feature of the innersurface of a mold employed in the present invention.

FIG. 4 is a cross-sectional view of an exemplary feature of the innersurface of another mold employed in the present invention.

FIG. 5 is a cross-sectional view of an exemplary feature of the innersurface of another mold employed in the present invention.

FIG. 6 is a cross-sectional view of an exemplary feature of the innersurface of another mold employed in the present invention.

FIG. 7 is a cross-sectional view of an exemplary feature of the innersurface of another mold employed in the present invention

FIG. 8 is a view showing a conventional casting apparatus employed inthe SC method.

FIG. 9 is a view showing a metallographic structure (cross section) ofan alloy ingot produced through a conventional SC method

DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

FIG. 1 is a photograph showing a cross section of an ingot (Nd: 30.0% bymass) of the present invention. The ingot of the present invention ischaracterized in that Nd-rich phase is uniformly dispersed in an almostdot-like manner. The dot-shaped, dispersed Nd-rich phase generally has amaximum width of 50 μm or less, and virtually no wire-shaped orrod-shaped Nd-rich phase, which is observed in conventional SCmaterials, is identified. In particular, Nd-rich phase which measures100 μm or more is substantially absent.

As used herein, the state “substantially absent” can be confirmedthrough observation of a cross section of an ingot in the followingmanner.

Specifically, a cross section of the ingot is polished, and arbitraryobservation areas on the cross section are observed under an SEM at amagnification of ×400. In each observation area, the rod-shaped Nd-richphase having a length of 100 μm or longer along a longitudinal directionis identified. When no Nd-rich phase of 100 μm or longer is observed inat least nine out of ten randomly selected observation areas, theNd-rich phase is evaluated as “substantially absent.”

The ingot of the present invention contains a minutely dispersed,dot-shaped Nd-rich phase, and in the SEM image, an area in which onlyNd-rich phase that measures 50 μm or less is dispersed accounts for atleast 50% the cross section. In other words, when ten arbitrary SEMphotographs (×400) are taken, at least five of them show no Nd-richphase measuring longer than 50 μm.

The ingot of the present invention is also characterized in that theingot contains reduced rod-shaped R-rich phase. In other words, asobserved on a cross section, R-rich phase having an aspect ratio of atleast 20 is “substantially absent.”

Regarding the above measurement, the state “substantially absent”generally refers to a level such that, among ten randomly selected SEM(×1,000) observation areas of a similarly polished cross section, thenumber of observation areas in which R-rich phase having an aspect ratioof at least 20 is present is about one or zero.

The ingot of the present invention is also characterized in that crystalgrains having a diameter of at least 1,000 μm as measured along themajor axis occupy an area percentage of at least 5%, leading toexcellent crystal orientation, and average R-rich phase spacing is 10 Enor less, leading to excellent sinterability of pulverized products.

The R-rich phase spacing is obtained through observation of a crosssection under an SEM. The spacing is an average of spacing values in adirection normal to the cast thickness direction obtained through imageprocessing or manual measurement of a photographic image.

In the ingot of the present invention, no substantial α-Fe is formeduntil the R content reaches approximately its stoichiometric value. Theexpression “no substantial α-Fe is formed” generally refers to a statein which, when presence of α-Fe is confirmed in ten arbitraryobservation areas of an arbitrary cross section of au ingot, α-Fe is notidentified in 90% or more of the observation areas. In theback-scattered electron image obtained through SEM, α-Fe is observed asdendrite-like dark portions.

The ingot of the present invention can be produced through the followingmethod FIG. 2 is an illustration showing one example apparatus forproducing a rare earth magnet alloy ingot of the present invention. Withreference to the illustration, the above method will be describedspecifically.

An outline of the centrifugal casting apparatus according to the presentinvention is shown in FIG. 2. FIG. 2 is one example of the presentinvention. The rare earth alloy containing alloys for R-T-B typemagnets, and the like, are melted in a melting chamber 1, for example,an alumina crucible 3, which is evacuated or contains an inert gasatmosphere for the purpose of the active properties thereof. The castingof the rare earth alloy is carried out by gradually tilting the crucibleso that the molten metal 31 of the rare earth alloy flows in the runner6, and is received by a cylindrical revolving body 5 having a base, forexample. By the rotation of the rotating body 5, the molten metal isdistributed from a plurality of holes 11 provided in the side surfacesof the rotating body 5 to the inner wall of the cylindrical mold 4 whichis on the outside of the rotating body. The cylindrical rotating body 5is manufactured so as to rotate around a rotation axis R which passesthrough the center of the circular base and is perpendicular to thebase. In addition, it is suitable for the rotating body to have thefunction of distributing the poured molten metal to the periphery, andin addition to a cylinder shape having a base, the rotating body mayhave any shape capable of distributing the molten metal such as a discshape, a cup shape angled outward toward the top, a cone shape angledoutward toward the bottom, and so on. However, a cylindrical shapehaving a plurality of holes 11 in the side walls as shown in the figureis preferable.

With this type of rotating body, when molten metal is poured into therotating body, the molten metal is dispersed to the periphery of therotating body due to the force of the rotation or the centrifugal force.In that situation, by reducing the thermal capacity of the rotatingbody, it is possible for the molten metal to be deposited and solidifyon the inner wall of the cylindrical mold without solidifying on therotating body.

In addition, in FIG. 2, the mold is a arranged horizontally. However, aslong as the position relationship with the rotating body is maintained,there is no problem in the cylindrical mold being arranged horizontally,or vertically, or being inclined.

In addition, by making an angle θ between the rotation axis R of therotating body 5 and the rotation axis L of the mold 4, it is possible toextend the deposition surface longitudinally along the entire surface ofthe mold, and thereby, it is possible to control the rate of depositionof the molten metal. By means of setting this angle θ, it is possible tospread the molten metal over a large surface area of the cylindricalmold, and as a result, it is possible to increase the rate ofsolidification. In addition, by making the rotation axis R of therotating body 5 variable, it is possible to spread the molten metal overan even greater surface area by varying the angle θ during casting.

For spreading the molten metal over the entire cylindrical mold 4, as analternative to having an angle θ between the rotation axis R of theabove-mentioned rotating body 5 and the rotation axis L of the mold 4,it is possible to obtain the same effects by moving the mold or therotating body back and forward in the direction of the rotation axis ofthe mold.

In addition, it is preferable for the rotating body and the mold to berotated in the same directions at different speeds of rotation. When therotating body and the mold are rotated in opposite directions, a splashphenomenon occurs readily in which when the molten metal strikes themold, it is scattered and does not stay on the mold, leading to areduction in the yield.

In addition, when the rotating body and the mold rotate in the samedirection with the same speed, the molten metal is deposited in a lineon the same surface of the mold, and it is not distributed over theentire surface of the mold. Consequently, the speed of rotation of thetwo should not be too close, and usually, the difference in the speed ofrotation of the two should be about 10% or greater and preferably 20% orgreater.

It is necessary to select the number of rotations of the rotating bodybased on the condition that the molten metal will strike the innersurface of the mold due to centrifugal force. The specific number ofrotations is determined in consideration of size of the rotating bodyand the mold, the direction in which the molten metal is sprayed outfrom the rotating body, the amount of molten metal which splasheswithout settling in the mold, and the like. When increasing the rate ofsolidification of the molten metal, it is preferable to determine thenumber of rotations such that the impact force of the molten metal onthe inner wall of the mold is greater.

In addition, it is necessary for the number of rotations of thecylindrical mold 4 to produce a centrifugal force of 1 G or greater suchthat the deposited and solidified alloy ingot 7 does not drop. Since thecooling effect is increased by the pressing of the molten metal onto theinner wall of the mold, a centrifugal force of 2 G or greater ispreferable.

The present invention is also characterized in that the inner surface ofthe rotatable mold 4 is rendered to be a non-smooth surface, therebyincreasing the cooling area of the mold, leading to an increase incooling performance and rate of cooling.

The non-smooth inner surface may have a curved surface as shown in FIG.3. However, grooves having a cross-section formed of linear segmentshaving angles therebetween as shown in FIGS. 4, 5, and 6 are preferable,since slippage of the molten metal from the mold surface, caused byshrinkage at solidification occurring at the moment of collision of themolten metal with the mold surface, can be prevented, thereby increasingadhesion of the molten metal with the mold and preventing deteriorationof heat transfer.

The depth of the grooves must be preset in consideration of volume andsurface area of the mold, specific heat, etc., and the depth isappropriately 0.5 nm to some nun. When the depth is considerably small,cooling effect becomes poor, resulting in failure to form desiredmetallographic structure, whereas when the depth is excessively large,removal of cast products from the inner surface becomes difficult.

The relationship between the size of molten metal droplets coming fromthe rotary body and the size and shape of the grooves is critical. Whenthe molten metal droplets are large and the grooves have a small widthand a large depth, a problem arises; i.e., molten metal droplets fail tocompletely penetrate the grooves, whereby a gap is formed between themold and deposited molten metal, possibly deteriorating the coolingeffect.

The mold is preferably formed of Cu, from the viewpoint of thermalconductivity. In addition to Cu, Fe is also employed without anyproblem.

When a conventional casting method is employed, R-rich phase iscrystallized along the formed columnar crystals and assumes the shape ofa rod. In addition, since the columnar crystals have a variety of growthdirections, R-rich phase is dispersed in a nonuniform manner. However,according to the present invention, the metallographic structure tendsto be formed of equiaxed crystals, since cooling performance of theinner surface of the mold is enhanced through increase in rate ofsolidification. In addition, R-rich phase is minutely crystallized andassumes virtually no rod shape. Thus, the R-rich phase is considered tobe dispersed more uniformly.

According to the casting method of the present invention, a molten metalis deposited in a mold and a subsequent molten metal is added to thedeposited molten metal that is being solidified. Since heat transfer forcooling is effected by mediation of a cast ingot, the thickness of theproduced ingot is limiter Generally, the maximum thickness is some tensof mm, and the preferred thickness is about 1 to about 10 mm. When thethickness is as thin as less than 1 mm, handling in the subsequentmagnet production step becomes cumbersome, whereas when the thickness isin excess of 10 mm, cooling performance of the surface of the ingotopposite the mold inner spice decreases.

Subsequently, an R-T-B type magnet alloy ingot which is produced throughthe casting method is pulverized, shaped, and sintered, to therebyproduce an anisotropic magnet of excellent characteristics.

Typically, pulverization is sequentially performed in the order ofhydrogen decrepitation, intermediate pulverization, andmicro-pulverization, to thereby produce a powder generally having a sizeof approximately 3 μm (FSSS).

In the present invention, hydrogen decrepitation includes a hydrogenabsorption step as a first step and a hydrogen desorption step as asecond step. In the hydrogen absorption step, hydrogen is caused to beabsorbed predominantly in the R-rich phase of alloy ingots in a hydrogengas atmosphere at 267 hPa to 50,000 hPa. The R-rich phase is expanded involume due to R hydride generated in this step, to thereby minutelyreduce the alloy ingots themselves or generate numerous micro-cracks.Hydrogen absorption is carried out within a temperature range of ambienttemperature to approximately 600° C. However, in order to increaseexpansion in volume of R-rich phase so as to effectively reduce theflakes in size, hydrogen absorption is preferably performed within atemperature range of ambient temperature to approximately 100° C. Thetime for hydrogen absorption is preferably one hour or longer. The Rhydride formed through the hydrogen absorption step is unstable in theatmosphere and readily oxidized. Thus, the hydrogen-absorbed product ispreferably subjected to hydrogen desorption treatment by maintaining theproduct at about 200 to about 600° C. in vacuum of 1.33 hPa or less.Through this treatment, R hydride can be transformed into a productstable in the atmosphere. The time for hydrogen desorption treatment ispreferably 30 minutes or longer. If the atmosphere is controlled forpreventing oxidation during steps to be carried out after hydrogenabsorption to sintering, hydrogen desorption treatment can also beomitted.

Alternatively, pulverization may be performed through intermediatepulverization and micro-pulverization without performing hydrogendecrepitation Intermediate pulverization is a pulverization step inwhich alloy flakes are pulverized in an inert gas atmosphere such asargon gas or nitrogen gas, to a particle size of, for example, 500 μm orless. Examples of pulverizers for performing this pulverization includea Brawn mill. In the present invention, if the alloy flakes have beensubjected to hydrogen decrepitation, the alloy flakes have already beenreduced minutely or have included numerous minute cracks generatedtherein. Thus, intermediate pulverization may be omitted.

Micro-pulverization is a pulverization step for attaining a particlesize of approximately 3 μm (FSSS). Examples of pulverizers forperforming the pulverization include a jet mill. Uponmicro-pulverization, the atmosphere is controlled to an inert gasatmosphere such as an argon gas atmosphere or nitrogen gas atmosphere.The inert gas may contain oxygen in an amount of 2% by mass or less,preferably 1% by mass or less. The presence of oxygen enhancespulverization efficiency and attains oxygen concentration of tee powderproduced through pulverization to 1,000 to 10,000 ppm, to therebyenhance oxidation resistance. In addition, abnormal gain growth duringsintering can be prevented.

In order to reduce friction between the powder and the inner wall of amold and to reduce friction generated among powder particles forenhancing orientation, a lubricant such as zinc stearate is preferablyadded to the powder during molding in magnetic field. The amount of thelubricant to be added is 0.01 to 1% by mass. Although the lubricant maybe added before or after micro-pulverization, the lubricant ispreferably mixed sufficiently, before molding in magnetic field, in aninert gas atmosphere such as argon gas or nitrogen gas by use of amixing apparatus such as a V-blender.

The powder pulverized to approximately 3 μm (FSSS) is press-molded inmagnetic field by use of a molding apparatus. The mold to be employed isfabricated from a magnetic material and a non-magnetic material incombination in consideration of the orientation of magnetic field in themold cavity. The pressure at molding is preferably 0.5 to 2 t/cm², andthe magnetic field in the mold cavity during molding is preferably 5 to20 kOe. The atmosphere during molding is preferably an inert gasatmosphere such as argon gas or nitrogen gas. However, if the powder hasbeen subjected to the aforementioned anti-oxidation treatment, moldingcan be performed in air.

Sintering is performed at 1,000-1,100° C. Prior to reaching thesintering temperature, a lubricant and hydrogen contained in themicro-powder must be completely removed from a compact to be sintered.The lubricant is removed by maintaining the compact preferably under theconditions: in vacuum of 1.33×10⁻² hPa or under an Ar flow atmosphere atreduced pressure; at 300 to 500° C.; and for 30 minutes or longer.Hydrogen is removed by maintaining the compact preferably under theconditions: in vacuum of 1.33×10⁻² hPa or less; at 700 to 900° C.; andfor 30 minutes or longer. The atmosphere during sintering is preferablyan argon gas atmosphere or a vacuum atmosphere of 1.3 3×10⁻² hPa orless. A retention time of one hour or longer is preferred.

After completion of sintering, in order to enhance the coercive force,the sintered product may be treated at 500 to 650° C. in accordance withneeds. An argon gas atmosphere or a vacuum atmosphere is preferred, anda retention time of 30 minutes or longer is preferred.

In addition to an R-T-B type magnet alloy, the casting method isapplicable to rare earth alloys such as misch metal-Ni alloys used foranodes of nickel-hydrogen batteries. According to quick-coolingsolidification involved in the method, segregation of Mn and othermetals can be prevented.

Another feature of the present invention lies in the provision of a filmhaving a thermal conductivity which is lower than the material of themold on the inner wall of the cylindrical mold which rotates. Thethermal conductivity of iron at normal temperatures is approximately 80W/mK. The thermal conductivity of the film provided on the inner surfaceof the mold is preferably 80 W/mK or less. Since this film acts as abarrier to the transmission to the mold of the heat of the molten metaldeposited on the mold, the temperature of the alloy ingot 7 deposited onthe inner surface of the mold at the initial stage of the casting ismaintained at a high temperature and does not fall very much. This hightemperature alloy ingot is pressed against the inner surface of the moldby the centrifugal force of the mold, and the surface which is incontact with the mold becomes as smooth as the inner surface of the moldwith almost no gaps between it and the mold. As a result, the heattransfer coefficient from the alloy ingot to the mold is converselyincreased, and molten metal which is deposited thereafter has anextremely fast rate of cooling.

According to the present invention, since the rate of cooling of themolten metal deposited on the mold is extremely fast, the particle sizeof the crystals of the alloy ingot for the R-T-B type ret is extremelyfine. It is possible to make the content of the alloy occupied bycrystals having a particle size of 10 μm or less be 50% or greater ofthe total alloy, preferably 70% or greater of the total alloy, and morepreferably 80% or greater of the total alloy. As a result, it ispossible for the cast ingot to display isotropically high magneticproperties even in its cast form.

In addition, in increasing the rate of cooling of the deposited moltenmetal, the rate of deposition of the molten metal on the mold is alsoimportant. In order to increase the rate of cooling, it is necessary toslow the rate of deposition, and preferably this rate is an average of0.1 mm/second or less, and more preferably an average of 0.05 mm/secondor less.

When the film is not provided on the inner wall surface of the mold, themolten metal deposited, at an initial casting stage, on the inner wallof the mold is rapidly cooled and solidified in such a form as providedat deposition. Thus, the surface (mold side) of the alloy ingot isimparted with a large number of irregularities, which considerablydeteriorate the rate of heat transfer from molten metal to besubsequently deposited onto the mold. As a result, crystal grain growthoccurs in the ingot, thereby generating a large number of crystal grainshaving a major grain size of at least 1,000 μm.

The crystal grain size of the alloy ingot may be determined throughobservation of a cross section of the alloy ingot in the followingmanner. Specifically, the cross section of the alloy ingot is polished,and an arbitrary observation area is observed under a polarizingmicroscope (×200) on the basis of the magnetic Kerr effect. The size ofeach crystal grain observed in the photographed observation area ismeasured through, for example, image processing. The ratio of the volumeof crystal grains having a specific grain size or less to the entirevolume of the alloy may be obtained in the following manner.Specifically, ten arbitrary observation areas of a cross section of analloy ingot are observed under the magnetic Kerr microscope (×200). Fromthe photographed ten observation areas, the total area corresponding tocrystal grains having the specific grain size or less is measuredthrough, for example, image processing, and the total area of suchcrystal grains is divided by the sum often photographed observationareas.

In the present invention, the method of attaching the film to the innerwall of the mold may be by any one of coating, plating, spraying, orwelding. For example, for coating, there are brush application,spraying, and the like, and for spraying, there are high-pressure-gasspraying, explosion spraying, plasma spraying, self-fluxing alloyspraying, and the like. In addition, for example, it is possible toadditionally provide a coated film on a spayed film. The thickness ofthe film is preferably in a range from 1 μm to 1 mm, and more preferably1 μm to 500 μm.

The material of the film is preferably selected from among a metal, aceramic, and a metal-ceramic composite. In addition it is preferable forthe film to comprises two or more layers which are made of differentmaterials. The material for the film is selected such that the formedfilm has thermal conductivity lower than that of the mold. Examples ofmetallic materials for forming the film include stainless steel, Ti, V,Cr, Mn, Be, Co, Ni, Cu, Nb, Mo, Ta, W, and alloys thereof. Even when themold and the film are both formed of Cu, the thermal conductivity of thefilm can be lowered as compared with the mold by, for example, providinga large number of micropores in the film. In a similar manner, Fe filmcan be used even when the mold is formed of Fe. Examples of ceramicmaterials for forming the film include boron nitride, sodium oxide, ironoxide, titanium oxide, aluminum oxide, calcium oxide, chromium oxide,zirconium oxide, tungsten oxide, vanadium oxide, barium oxide, manganeseoxide, magnesium oxide, silicon oxide, rare earth oxides, tungstencarbide, chromium carbide, niobium carbide, titanium carbide, andcomposite ceramics thereof. In addition, composite films formed of theaforementioned metallic materials and ceramics may also be employed.

The inner wall surface of the rotatable mold may be rendered to be anon-smooth surface, thereby increasing the cooling area of the mold,whereby a film is provided thereon, leading to an increase in coolingperformance and rate of cooling. The non-smooth inner wall surface mayhave a curved surface. However, grooves having a cross-section formed oflinear segments having angles therebetween are preferable, sinceslippage of the molten metal from the mold surface, which wouldotherwise be caused by shrinkage at solidification occurring at themoment of collision of the molten metal with the mold surface, can beprevented, thereby increasing adhesion of the molten metal with the moldand preventing deterioration of heat transfer. The depth of the groovesmust be preset in consideration of volume and surface area of the mold,specific heat, etc., and the depth is appropriately 0.5 mm to some mm.When the depth is considerably small, cooling effect becomes poor,resulting in failure to form desired metallographic structure, whereaswhen the depth is excessively large, removal of cast products from theinner surface becomes difficult.

According to the present invention, a new molten metal is cast on thesufficiently cooled ingot, and this step is repeated, to thereby producea thickness-increased alloy ingot having a metallographic structureincluding microcrystal grains. In practice, the alloy ingot preferablyhas a thickness of at least 1 mm, more preferably at least 5 mm, mostpreferably at least 10 mm.

When the alloy ingot of the present invention is hot worked at hightemperature in an inert gas atmosphere or in vacuum, anisotropy can beprovided Examples of methods of preferable hot workings includedie-upsetting, rolling, forging, and pressing. The temperature ofdeformation is preferably 500 to 1,100° C., more preferably 600 to 800°C. The preferable pressure is at least 0.5 ton/cm², more preferably atleast 1 ton/cm².

In addition, it is possible to increase the coercive force and thesquareness properties of the alloy ingot of the present invention bycarrying out a heat treatment at 400 to 1,000° C. in a vacuum or in aninactive gas atmosphere after the casting or after the hot workingprocess. Alternatively, it is possible to magnetize and further increasethe coercive force by carrying out, after the casting or after the hotworking process, a heat treatment at 1,000 to 1,100° C. in a vacuum oran inactive gas atmosphere, and then carrying out a heat treatment at400 to 1,000° C. in a vacuum or in an inactive gas atmosphere.

The R-T-B type magnetic alloy of the present invention comprises atleast one element selected from among Nd, Pr, and Dy in a total amountof 11.8 to 16.5% by atom, and B in an amount of 5.6 to 9.1% by atom,with a balance being T (T represents a substance predominantlycomprising Fe, with a portion of Fe atoms being optionally substitutedby Co, Ni, Cu, Al, Ga, Cr, and Mn.). When the total amount of at leastone element selected from among Nd, Pr, and Dy is less than 11.8% byatom, R-rich phase content of the alloy is poor, thereby deterioratingmagnetic characteristics, whereas when the total amount is in excess of16.5% by atom, non-magnetic R-rich phase content increases excessively,thereby deteriorating magnetization. When the amount of B is less than5.6% by atom, magnetic characteristics are deteriorated because of aninsufficient B content, whereas when the amount of B is in excess of9.1% by atom, nonmagnetic B-rich phase (R₁₊₆T₄B₄ phase) contentincreases excessively, thereby deteriorating magnetization. Therefore,the alloy is formed of at least one element selected from among Nd, Pr,and Dy in a total amount of 11.8 to 16.5% by atom and B in an amount of5.6 to 9.1% by atom, with a balance being T.

In order to generate further minute crystal grains in the R-T-B typemagnetic alloy, high-melting-point metals such as Ti, V, Cr, Mn, Zr, Nb,Mo, Hf, Ta, and W may be added. In this case, the total amount of themetals is preferably regulated to 1% by mass or less so as to preventdeterioration in magnetization.

The R-T-B type magnetic alloy of the present invention exhibitsexcellent magnetic characteristics isotropically even when in an ingotstate after casting. Therefore, isotropic magnets exhibiting excellentmagnetic characteristics can be produced through mere mechanicalprocessing; e.g., cutting, grinding, polishing, or blanking, of an alloyingot to form a predetermined shape. Briefly, steps of pulverization,pressing in a magnetic field, and sintering, which are necessary forproducing the conventional sintered magnets, or steps of pulverizationand press-molding or injection molding performed in production of bondedmagnets can be omitted, thereby reducing production costs. Since themagnets produced through mechanical processing have a density higherthan that of bonded magnets, strong magnets exhibiting highmagnetization can be produced.

On the basis of deposition of a molten metal on the inner wall of thecylindrical mold, magnets of a cylindrical shape can be produceddirectly from a molten metal. In this case, in view of limitations ofproduction apparatuses, the outer diameter of the cylindrical R-T-B typemagnets is preferably controlled to 100 mm or more.

The R-T-B type magnet of the present invention contains a rare earthcomponent and Fe, which are highly prone to oxidization Thus, the magnetis preferably coated with a resin or a metal such as Ni or Al. Morepreferably, the magnet is sequentially coated with a resin and a metal.

The aforementioned R-T-B type alloy ingot of the present inventioncontains crystal grains, most of which are minute. Thus, even when thealloy ingot is pulverized, magnetic characteristics are not greatlydeteriorated. The ingot is pulverized to a particle size of 500 μm orless, and the powder is mixed with an epoxy resin or a similar resin andthe resultant mixture press-molded; or the powder is mixed with Nylon ora similar resin and the resultant Mature injection-molded, to therebyproduce a bonded magnet. The production method for R-T-B type magnetalloy ingot of the present invention attains productivity higher thanthat attained by the aforementioned ultra-rapid-cooling productionmethod, thereby providing low-price R-T-B type bonded magnet alloypowder.

According to the present invention, casting is continued while rareearth alloy deposited on the inner wall surface of the mold are scrapedby use of a scraper or a similar device, thereby producing flat rareearth alloy flakes having a maximum length of 5 cm and a thickness of 1mm or less. After completion of casting, the flakes are heat-treated at400 to 1,000° C. in vacuum or in an inert gas atmosphere, to therebyenhance coercive force and squareness properties. Alternatively, theflakes are heat-treated at 1,000 to 1,100° C. in vacuum or in an inertgas atmosphere and, subsequently, at 400 to 1,000° C. in vacuum or in aninert gas atmosphere, to thereby further enhance magnetization andcoercive force.

The R-T-B type alloy flakes contain crystal grains, most of which areminute. Thus, even when the alloy flakes are pulverized, magneticcharacteristics are not greatly deteriorated. The flakes are pulverizedto a particle size of 500 μm or less, and the powder is mixed with anepoxy resin or a similar resin and the resultant mixture press-molded;or the powder is mixed with Nylon or a similar resin and the resultantmixture injection-molded, to thereby produce a bonded magnet. If thebonded magnets are produced by pulverizing alloy flakes is preferablebecause low-price alloy powder can be provided by virtue of highpulverization efficiency, as compared with pulverization of alloy ingotsand obtaining the bonded magnets.

The production method for rare earth alloy of the present invention canprovide an R-T-B TYPE exchange spring magnet alloy ingot comprising atleast one element selected from among Nd, Pr, and Dy in a total amountof 1 to 12% by atom and B in an amount of 3 to 30% by atom, with abalance being T (T represents a substance predominantly comprising Fe,with a portion of Fe atoms being optionally substituted by Co, Ni, Cu,Al, Ga, Cr, and Mn.); containing crystal grains of the hard magneticphase and the soft magnetic phase having a crystal grain size of 1 μm orless in a total volume of at least 50%, preferably at least 70%, morepreferably at least 80%, on the basis of the entire volume of the alloy.The thickness of the alloy ingot is 1 mm or greater, preferably 5 mm orgreater, and even more preferably 10 mm or greater.

The R-T-B type exchange spring magnet of the present invention containsa hard magnetic phase formed of R₂T₁₄B, exhibiting large anisotropicmagnetic field, and a soft magnetic phase formed of at least one speciesselected from among Fe, Fe, and Fe₃B, exhibiting high saturationmagnetization.

In the R-T-B type exchange spring magnet of the present invention, inorder to generate further minute crystal grains, high-melting-pointmetals such as Vt, V, Cr, Mn, Zr, Nb, Mo, H, Ta, and W may be added. Inthis case, the total amount of the metals is preferably regulated to 1%by weight or less so as to prevent deterioration in magnetization.

When the R-T-B type exchange spring magnet is produced by thecentrifugal casting method of the present invention, the rate ofrotation of the mold is preferably regulated to generate centrifugalforce of at least 2 G, more preferably at least 5 G, most preferably atleast 10 G, such that the cooling effect is increased by bringing themolten metal into contact with the inner wall of the mold.

The rate of rotation of the rotary body is preferably regulated toimpart, to the molten metal, centrifugal force of at least 5 G, morepreferably at least 20 G, most preferably at least 30 G, such thatcollision of the molten metal with the inner wall surface of the mold isintensified, to thereby increase the solidification rate of the moltenmetal.

Rate of deposition of a molten metal on the inner wall surface of themold is also critical. The average deposition rate for increasing thecooling rate of the deposited molten metal is 0.1 mm/sec or less,preferably 0.05 mm/sec or less, more preferably 0.03 mm/sec or less.

After casting, the R-T-B type exchange spring magnet alloy ingot isheat-treated at 400 to 1,000° C. in vacuum or in an inert gasatmosphere, to thereby enhance coercive force and squareness properties.

The exchange spring magnet alloy ingot of the present invention exhibitslarge percent spring-back, i.e., even when magnetization decreases in areverse magnetic field, magnetization is almost completely restored tothe initial value through controlling the magnetic field to 0. Inaddition, the alloy also exhibits excellent magnetic characteristicsisotropically at an ingot state. Therefore, exchange spring magnetsexhibiting excellent magnetic characteristics isotropically can beproduced only trough mechanical processing; e.g., cutting, grinding,polishing, or blanking, of an alloy ingot to form a predetermined shape.Briefly, steps of pulverization and press molding or injection molding,which are required for producing the conventional bonded magnets can beomitted, thereby reducing production costs. In addition, since themagnets produced through mechanical processing have a density higherthan that of bonded magnets, strong magnets exhibiting highmagnetization can be produced.

On the basis of deposition of a molten metal on the inner wall of thecylindrical mold, the R-T-B type exchange spring magnets of acylindrical shape can be produced directly from a molten metal. In thiscase, in view of limitations of production apparatuses, the outerdiameter of the cylindrical magnets is preferably controlled to 100 mmor more.

When the exchange spring magnet alloy ingot of the present invention isintentionally deformed at high temperature in an inert gas atmosphere orin vacuum, anisotropy can be provided. Examples of methods ofdeformation include die-upsetting, rolling, forging, and pressing. Thetemperature of deformation is preferably 400 to 1,000° C., morepreferably 600 to 800° C. The pressure required for deformation is atleast 0.5 ton/cm², more preferably at least 1 ton/cm².

The exchange spring magnet of the present invention contains a rareearth component and Fe, which are highly prone to oxidization. Thus, themagnet is preferably coated with a resin or a metal such as Ni or Al.More preferably, the magnet is sequentially coated with a resin and ametal.

The aforementioned exchange spring magnet alloy ingot of the presentinvention contains crystal grains, most of which are minute. Thus, evenwhen the alloy ingot is pulverized, magnetic characteristics are notgreatly deteriorated. The ingot is pulverized to a particle size of 500μm or less, and the powder is mixed with an epoxy resin or a similarresin and the resultant mixture press-molded; or the powder is mixedwith Nylon or a similar resin and the resultant mixtureinjection-molded, to thereby produce a bonded magnet. When the exchangespring net alloy ingot is produced by the production method of thepresent invention, the alloy attains productivity higher than thatattained by the aforementioned ultra-rapid-cooling production method,thereby providing low-price alloy powder.

EXAMPLES

Below, the Examples of the present invention and the ComparativeExamples will be explained.

Example 1

Elemental neodymium, ferroboron, cobalt, aluminum, copper, and iron weremixed so as to obtain the following composition: Nd: 30.0% by mass: B:1.00% by mass; Co: 1.0% by mass; Al: 0.30% by mass; and Cu: 0.10% bymass, a balance being iron. The resulting mixture was melted in analumina crucible in an argon gas atmosphere at 1 atm by use of ahigh-frequency induction melting furnace. The resulting molten mixturewas subjected to casting by use of the apparatus shown in FIG. 2.

The mold has an inside diameter of 500 mm and a length of 500 mm, andgrooves shown in FIG. 7 having a depth of 1 mm and a bottom width of 5mm are provided with intervals of 3 mm in the inner Race of the mold.

The rotary receptacle has an inside diameter of 250 mm and eight holesof 2 mm in diameter formed in the surrounding wall of the receptacle.

The angle θ formed by the axis of rotation of the rotary receptacle andthe axis of rotation of the mold was fixed to 25′, and the averagemolten metal deposition rate for deposition on the inner wall of themold was 0.01 cm/sec.

The rotational speed of the mold was set to 189 rpm so as to generate acentrifugal force of 10 G. The rotational speed of the rotary receptaclewas 535 rpm so as to impose a centrifugal force of about 40 G on moltenmetal.

The thus-obtained alloy ingot had a thickness of 6 to 8 mm as measuredat a central portion of the cylindrical mold and a thickness of 11 to 13mm as measured at thickest portions located in the vicinity of oppositeend portions. The micro-structure of the cross section was observed as aback-scattered electron image under an electron microscope. The resultsare shown in Table 1.

TABLE 1 R-rich phase Crystal grain size Composition (% by mass) Aspectratio Spacing (major axis) Nd Cu B Co Al Fe ≦50 μm of ≧20 μm of ≧1,000μm Remarks Ex. 1 30.0 0.10 1.00 1.0 0.30 bal. 60% Not found 6 20% Mold:rough surface, 10 G Ex. 2 28.0 0.10 1.00 1.0 0.30 bal. 65% Not found 310% Mold: smooth, surface, 15 G Comp. Ex. 1 30.0 0.10 1.00 1.0 0.30 bal.10% Moderate 10 50% Mold: smooth surface, 2.5 G Comp. Ex. 2 30.0 0.101.00 1.0 0.30 bal. 0 Many 4 0 SC method

Example 2

Elemental neodymium, ferroboron, cobalt, aluminum, copper, and iron weremixed so as to obtain the following composition: Nd: 28.0% by mass; B:1.00% by mass; Co: 1.0% by mass; Al: 0.30% by mass; and Cu: 0.10% bymass, a balance being iron. The resulting mixture was melted in analumina crucible in an argon gas atmosphere at 1 atm by use of ahigh-frequency induction melting furnace. The resulting molten mixturewas subjected to casting by use of the apparatus shown in FIG. 2.

A mold and a rotary receptacle having the same dimensions as those ofExample 1 were employed. However, the inner surface of the mold wassmooth, and the rotational speed of the mold was set to 231 rpm so as togenerate a centrifugal force of 15 G.

The same conditions as described in relation to the Example wereemployed for the rotary receptacle.

The results are shown in the above Table 1.

Comparative Example 1

A mixture having a composition similar that of the alloy of Example 1was prepared, and the mixture was melted in a manner similar to that ofExample 1 and cast by use of a casting apparatus similar to thatemployed in Example 1. However, the inner surface of the mold had nogrooves, and the surface was polished to be smooth in advance by use ofsand paper (No. 240). The rotational speed of the mold was adjusted soas to generate a centrifugal force of 2.5 G The alloy ingot obtainedthrough the above casting had a thickness of 7 to 8 mm as measured at acentral portion of the cylindrical mold and a thickness of 12 to 13 mmas measured at thickest portions located in the vicinity of opposite endportions. Its a manner similar to that employed in Example 1, themicro-structure of the cross section was observed as a back-scatteredelectron image. The results are shown in Table 1.

Comparative Example 2

A mixture having a composition similar that of the alloy of Example 1was prepared, and the mixture was melted under argon at 1 atm and castby use of a casting apparatus shown in FIG. 8 for the SC method. Thewater-cooling copper roller 23 had an outer diameter of 400 mm, and theroller was rotated at a peripheral velocity of 1 m/s, thereby yielding aflaky alloy ingot having a mean thickness of 0.32 mm. Themicro-structure of the cross section of the thus-obtained alloy ingotwas observed as a back-scattered electron image. The results are shownin Table 1.

Example 3

Example 3 describes a production example of sintered magnets. The alloyflakes produced in Example 1 were pulverized in the order of hydrogendecrepitation, intermediate pulverization, and micro-pulverization.Hydrogen absorption step—the first step of hydrogen decrepitation—wasperformed under the conditions: 100% hydrogen atmosphere, atmosphericpressure, and retention time of 1 hour. The temperature of the metalflakes at the start of hydrogen absorption reaction was 25° C. Hydrogendesorption step—subsequent step—was performed under the conditions:vacuum of 0-13 hPa, 500° C., and retention time of 1 hour. Intermediatepulverization was performed by use of a Brawn mill, and thehydrogen-decrepitated powder was pulverized in a 100% nitrogenatmosphere to a particle size of 42.5 μm or less. To the resultantpowder, zinc stearate powder was added in an amount of 0.07% by mass.The mixture was sufficiently mixed in a 100% nitrogen atmosphere by useof a V-blender, and then micro-pulverized by use of a jet mill in anitrogen atmosphere incorporated with oxygen (4,000 ppm), to a particlesize of 3.2 μm (FSSS). The resultant powder was sufficiently mixed againin a 100% nitrogen atmosphere by use of a V-blender. The obtained powderwas found to have an oxygen concentration of 2,500 ppm. Through analysisof the carbon concentration of the powder, the zinc stearate content ofthe powder was calculated to be 0.05% by mass.

Subsequently, the thus-obtained powder was press-molded in a 100%nitrogen atmosphere and a lateral magnetic field by use of a moldingapparatus. The molding pressure was 1.2 t/cm², and the magnetic field inthe mold cavity was controlled to 15 kOe.

The thus-obtained compact was maintained sequentially in vacuum of1.33×10⁻⁵ hPa at 500° C. for one hour and in vacuum of 1.33×10⁻⁵ hPa at800° C. for two hours, and was further maintained in vacuum of 1.33×10⁻⁵hPa at 1,060° C. for two hours for sintering. The density of thesintered product was as sufficiently high as 7.5 g/cm³ or more. Thesintered product was further heat-treated at 540° C. for one hour in anargon atmosphere.

Magnetic characteristics of the sintered product were measured by meansof a direct-current BH curve tracer, and the results are shown in Table2. In addition, a cross-section of the sintered product wasmirror-polished, and the polished surface was observed under apolarizing microscope. The obtained mean crystal m size was 15 to 20 μm,with substantial homogeneity in size.

Comparative Examples 3 and 4

In Comparative Examples 3 and 4, each of the alloy flakes produced inComparative Examples 1 and 2 was pulverized in a manner similar to thatof Example 3, to thereby produce a powder having a particle size of 3.3μm (FSSS). The powder was found to have an oxygen concentration of 2,600ppm. In a manner similar to that of Example 3, each powder was molded ina magnetic field and sintered, to thereby produce an anisotropic magnet.Magnetic characteristics of the thus-produced sintered magnets are shownin Table 2.

TABLE 2 Br iHc BH_(max) Remarks Ex. 3 14.22 kG 11.38 kOe 47.8 MGOe Mold:rough surface, 10 G Comp. Ex. 3 14.18 kG  9.32 kOe 47.3 MGOe Mold:smooth surface, 2.5 G Comp. Ex. 4 14.07 kG 11.42 kOe 46.8 MGOe SC method

Table 2 indicates that the magnet produced in Comparative Example 3exhibits a coercive force (iHc) lower, by 2 kOe or more, t that of themagnet produced in Example 3, conceivably due to a poor dispersion stateof R-rich phase. The magnet produced in Comparative Example 4 exhibits aBr as low as 0.15 kG conceivably due to crystal orientation inferior tothat of the alloy of the present invention.

Example 4

Elemental neodymium elemental praseodymium, ferroboron, aluminum,eletrolytic copper, electrolytic cobalt, and electrolytic iron weremixed so as to obtain the following composition: Nd: 10.4% by atom(23.0% by mass); Pr: 3.2% by atom (7.0% by mass); B: 6.0% by atom (1.0%by mass); Al: 0.7% by atom (0.30% by mass); Cu: 0.1% by atom (0.10% bymass); Co: 1.1% by atom (1.0% by mass); and a balance of iron. Theresulting mixture was melted in an alumina crucible in an argon gasatmosphere through high-frequency induction melting. The resultingmolten mixture was subjected to casting under the following conditionsusing the device shown in FIG. 2.

The cylindrical mold which is made of iron (thermal conductivity at 27°C.: 80.3 W/mK) has an inside diameter of 500 mm and a length of 500 mm.Coating film (thermal conductivity at 27° C.: 12.6 W/mK) having acomposition of Ni: 80% by mass and Cr: 20% by mass, and a thickness of100 μm was provided on the inner wall of the mold through plasmaspraying. The rotary body was a cylindrical receptacle having an insidediameter of 250 mm and eight holes of 3 mm in diameter formed in thesurrounding wall of the receptacle. The axis L of rotation of thecylindrical mold is set to the horizontal direction. The angle θ ofinclination formed by the axis R of rotation of the receptacle and theaxis L of rotation of the cylindrical mold was fixed to 25 G duringmolding, and the average molten metal deposition rate for deposition onthe inner wall of the mold was adjusted to 0.05 mw/sec. The rotationalspeed of the mold was controlled so as to generate a centrifugalacceleration of 10 G. The rotational speed of the rotary receptacle wascontrolled so as to impose a centrifugal force of about 20 G on moltenmetal.

The thus-obtained alloy ingot had a thickness of 8 mm as measured at acentral portion of the cylindrical mold and about 10 mm as measured atthickest portions located in the vicinity of opposite end portions. Across section of the alloy ingot was observed under a polar microscopefor determining the crystal grain size. Through observation, the percentarea in which the crystal grain size of 10 μm or less occupied was foundto be 95%.

Then, a cube (side: 7 mm) was cut from the alloy ingot, and magneticcharacteristics of the cube were determined by means of a BH curvetracer. The alloy was found to exhibit the following characteristics:residual magnetic flux density Br=8.6 kG; coercive force iHc=10.2 kOe;and maximum energy product (BH)_(max)=14.2 MGOe, and thesecharacteristics were almost equivalent in three axial directions. Theresults indicate that the alloy suitably provides an alloy ingot forisotropic magnets.

Example 5

The same materials were mixed so as to obtain the composition similar tothat of the alloy produced in Example 4. The resulting mixture wasmelted in an alumina crucible in an argon gas atmosphere throughhigh-frequency induction melting. The resulting molten mixture wassubjected to casing by use of the apparatus and under the conditionssimilar to those employed in Example 4. However, instead of the abovecoating film, boron nitride (BN) film (thermal conductivity at 27° C.:17 to 42 W/mK) was provided in a thickness of 10 μm through spraycoating on the inner wall surface of the mold.

A cross section of the thus-obtained alloy ingot was observed under apolarizing microscope for determining the crystal grain size. Throughobservation, the percent area in which the crystal grain size of 10 μmor less occupied was found to be 88%.

Then, a cube (side: 7 mm) was cut from the alloy ingot, and magneticcharacteristics of the cube were determined by means of a BH curvetracer. The alloy was found to exhibit the following characteristics:Br=8.6 kG; iHc=10.1 kOe; and (BH)_(max)=14.0 MGOe, and thesecharacteristics were almost equivalent in three axial directions. Theresults indicate that the alloy suitably provides an alloy ingot forisotropic magnets.

Example 6

Elemental neodymium, ferroboron, and electrolytic iron were mixed so asto obtain the following composition: Nd: 4.6% by atom (12.5% by mass);B: 15.2% by atom (3.1% by mass); and a balance of iron. The resultingmixture was melted in an alumina crucible in an argon gas atmospherethrough high-frequency induction melting. The resulting molten mixturewas subjected to casting by use of the apparatus shown in FIG. 2 andunder the conditions described below.

The cylindrical mold which is made of iron (thermal conductivity at 27°C.: 80.3 W/mK) has an inside diameter of 500 mm and a length of 500 mm.Coating film (thermal conductivity at 27° C.: 12.6 W/mK) having acomposition of Ni: 80% by weight and Cr: 20% by weight, and a thicknessof 500 μm was provided on the inner wall of the mold through plasmaspraying. The rotary body was a cylindrical receptacle having an insidediameter of 250 mm and eight holes of 2 mm in diameter formed in thesurrounding wall of the receptacle. The axis L of rotation of thecylindrical mold is set to the horizontal direction. The angle θ ofinclination formed by the axis R of rotation of the rotary receptacleand the axis L of rotation of the cylindrical mold was fixed to 25°during molding.

The average molten metal deposition rate for deposition on the innerwall of the mold was adjusted to 0.02 mm/sec. The rotational speed ofthe mold was controlled so as to generate a centrifugal acceleration of20 G. The rotational speed of the rotary receptacle was controlled so asto impose a centrifugal force of about 40 G on molten metal.

A cross section of the thus-obtained alloy ingot was observed under apolarizing microscope for determining the crystal in size. Throughobservation, the percent area in which the crystal grain size of 1 μm orless occupied was found to be 65%.

Then, a cube (side: 7 mm) was cut from the alloy ingot, and magneticcharacteristics of the cube were determined by means of a BH curvetracer. The alloy was found to exhibit the following characteristics:Br=11.8 kG; iHc=3.0 kOe; and (BH)_(max)=14.9 MGOe, and thesecharacteristics were almost equivalent in three axial directions. Theresults indicate that the alloy suitably provides an alloy ingot forisotropic magnets. To the magnet obtained from the ingot, a reversemagnetic field of 2.5 kOe was applied after magnetization. When theapplication of the magnetic field was stopped (i.e., to 0 kG), Br wasrestored to 95% the initial value; i.e., remarkable sprung back.Therefore, the magnet was identified to be an isotropic exchange springmagnet.

Comparative Example 5

The same materials were mixed so as to obtain the composition similar tothat of the alloy produced in Example 4. The resulting mixture wasmelted in an alumina crucible in an argon gas atmosphere throughhigh-frequency induction melting. The resulting molten mixture wassubjected to casting by use of the apparatus and under the conditionssimilar to those employed in Example 4. However, no coating film wasprovided on the inner wall surface of the mold, and the molten alloy wasdeposited and solidified directly on the inner wall surface of the ironmold A cross section of the thus-obtained alloy ingot was observed undera polarizing microscope for determining the crystal grain size. Throughobservation, a large number of columnar crystals having a major size of1 mm or more were observed, and the percent area in which the crystalgrain size of 10 μm or less occupied was found to be as low as 3%.

Then, a cube (side: 7 mm) was cut from the alloy ingot, and magneticcharacteristics of the cube were determined by means of a BH curvetracer. The alloy was found to exhibit the following characteristics:Br=3.0 kG; iHc=0.8 kOe; and (BH)_(max)=0.4 MGOe. These characteristicsexhibited their highest when measured in a plane normal to the moldsurface, but are considerably low as compared with those of Example 4.

Comparative Example 6

The same materials were mixed so as to obtain the composition similar tothat of the alloy produced in Example 6. The resulting mixture wasmelted in an alumina crucible in an argon gas atmosphere throughhigh-frequency induction melting. The resulting molten mixture wassubjected to casting by use of the apparatus and under the conditionssimilar to those employed in Example 4. However, no coating film wasprovided on the inner wall surface of the mold, and the molten alloy wasdeposited and solidified directly on the inner wall surface of the ironmold. A cross section of the thus-obtained alloy ingot was observedunder a polarizing microscope for determining the crystal grain size.Through observation, a large number of columnar crystals having a majorsize of 1 mm or more were observed. However, many portions of dendriticphase in which no magnetic domain was identified were observed, and thephase seemed to prevent growth of columnar crystals. On the basis of aback-scattered electron image captured by a scanning electron microscopeand by means of an energy dispersion X-ray analyzer, the dendritic phasein which no magnetic domain was observed was identified to α-Fe. Inaddition, through observation of the alloy ingot under a polarizingmicroscope, the percent area in which the crystal grain size of 10 μm orless occupied was found to be as low as 3%.

Then, a cube (side: 7 mm) was cut from the alloy ingot, and magneticcharacteristics of the cube were determined by means of a BH curvetracer. The alloy was found to exhibit the following characteristics:Br=1.8 kG; iHc=0.2 kOe; and (BH)_(max)=unmeasurable. Thesecharacteristics exhibited their highest when measured in a plane normalto the mold surface, but are considerably low as compared with those ofExample 6.

Example 7

The alloy ingot of Example 4 was pulverized to 500 μm or less in anargon gas atmosphere by use of a stamp mill, and iHc of the resultantpowder was found to be 9.5 kOe as measured by means of a vibratingsample magnetometer (VSM), indicating a small decrease in iHc. The alloypowder was mixed with epoxy resin (3% by weight), and the resultantmixture was press-formed at 6 ton/cm² in an argon gas atmosphere. Theresultant compact was fired at 180° C. in an argon gas atmosphere, tothereby cure the epoxy resin. After completion of curing, the density ofthe product was found to be 5.8 g/cm³. Magnetic characteristics of theproduct were determined by means of a BH curve tracer. The product wasfound to exhibit the following characteristics: Br=6.6 kG; coerciveforce iHc=9.1 kOe; and maximum energy product (BH)_(max)=8.4 MGOe.

Example 8

The alloy ingot of Example 4 was heated at 550° C. in vacuum for onehour. Then, a cube (side: 7 mm) was cut from the heat-treated alloyingot, and magnetic characteristics of the cube were determined by meansof a BH curve tracer. The alloy was found to exhibit the followingcharacteristics: Br=8.7 kG; iHc=11.2 kOe; and (BM)_(max)=14.9 MGOe, andthese characteristics were almost equivalent in three axial directions.

Example 9

The alloy ingot of Example 8 was pulverized to 500 μm or less in anargon gas atmosphere by use of a stamp mill, and iHc of the resultantpowder was found to be 10.5 kOe as measured by means of a VSM,indicating a small decrease in iHc. By use of the alloy powder and in amanner similar to that of Example 7, a bonded magnet having a density of5.8 g/cm³ was produced. Magnetic characteristics of the bonded magnetwere determined by means of a BH curve tracer. The magnet was found toexhibit the following characteristics: Br=6.8 kG; iHc=10.2 kOe; and(BH)_(max)=8.9 MGOe.

Example 10

The alloy ingot of Example 4 was heated at 1,020° C. in an argonatmosphere for two hours, followed by an additional heat treatment at550° C. in vacuum for one hour. Then, a cube (side: 7 mm) was cut fromthe heat-treated alloy ingot, and magnetic characteristics of the cubewere determined by means of a BH curve tracer. The alloy was found toexhibit the following characteristics: Br=8.9 kG; iHc=11.3 kOe; and(BH)_(max)=15.5 MGOe. These characteristics were almost equivalent inthree axial directions.

Example 11

The alloy ingot of Example 10 was pulverized to 500 μm or less in anargon gas atmosphere by use of a stamp mill, and iHc of the resultantpowder was found to be 10.7 kOe as measured by means of a VSM,indicating a small decrease in iHc. By use of the alloy powder and in amanner similar to that of Example 7, a bonded magnet having a density of5.8 g/cm³ was produced. Magnetic characteristics of the bonded magnetwere determined by means of a BH curve tracer. The magnet was found toexhibit the following characteristics: Br=6.9 kG; iHc=14 kOe; and(BH)_(max)=9.3 MGOe.

Comparative Example 7

The alloy ingot of Comparative Example 4 was pulverized to 500 μm orless in an argon gas atmosphere by use of a stamp mill, and iHc of theresultant powder was found to be as low as 0.4 kOe as measured by meansof a vibrating sample magnetometer (VSM). By use of the alloy powder andin a manner similar to that of Example 7, a bonded magnet having adensity of 5.8 g/cm³ was produced. Magnetic characteristics of thebonded magnet were determined by means of a BR curve tracer. The magnetwas found to exhibit the following characteristics: Br_(max)=2.3 kG;iHc=0.3 kOe; and (BH)_(max)=0.1 MGOe, which were considerablyunsatisfactory.

Example 12

In order that the composition comprise Nd: 14.7% by atom (32.0% bymass), B: 6.1% by atom (1.0% by mass), Al: 0.7% by atom (0.30% by mass),Cu: 1.0% by atom (1.0% by mass), Nb: 0.4% by atom (0.5% by mass), andthe remainder iron, each of the starting materials of metal neodymium,ferro-boron, aluminum, electrolytic copper, ferro-niobium, electrolyticiron were mixed, and melted by high frequency induction heating using analumina crucible in an argon gas atmosphere, and then casting wascarried out using the same apparatus and conditions as were used inExample 5. Thereafter, without removing the deposited alloy ingot fromthe cylindrical mold, onto this alloy ingot, an alloy ingot having thesame composition as the previous composition was deposited under thesame manufacturing conditions.

The thickness of the obtained alloy ingot was 16 mm at the centersection of the cylindrical mold, and approximately 20 mm at the thickestsections in the vicinity of both edges. As a result of the measurementof the diameter of the crystals of the alloy ingot using a polarizingmicroscope, it was found that the area occupied by crystals having aparticle size of 10 μm or less was 83%.

A cube have sides of 7 mm was cut out from the portion of the alloywhich was deposited and solidified later, and the magnetic propertieswere measured using a BH curve tracer. The properties in threedirections were approximately the same and Br=8.2 kG, iHc=4 kOe, and(BH) max=12.9 MGOe.

Example 13

A portion having a thickness of 16 to 18 mm was cut out from the alloyingot of Example 12, and was enclosed in an evacuated iron vessel havinga thickness of 3.2 mm. The iron vessel in which his alloy ingot wasenclosed was placed in an atmospheric oven set at 800° C., sufficientlyheated, and then rolled by passing between rollers set to have a draftof 30%. Thereafter, the vessel was returned to the atmospheric ovenmaintained at 800° C., sufficiently heated, and then rolled again withthe interval between the rollers reduced with a draft of 30%. Thisrolling operation was repeated for a total of 4 times, and the alloyingot was rolled to a thickness of 4.0 mm. Two pieces of this alloyingot were laminated together and processed to give a cube having sidesof 7 mm. The magnetic properties were measured by a BH curve trace, andit was found that the magnetic force in the direction of pressing wasstrongest, and the magnetic properties in this direction were Br=12.0kG; iHc=12.9 kOe, and (BH) max=28.7 MGOe.

Example 14

The alloy ingot which was hot rolled in Example 13 was heat treated for1 hour at 550° C. in a vacuum. Thereafter, two pieces of this alloyingot were laminated together and processed to give a cube having sidesof 7 mm. The magnetic properties were mead by a BH curve tracer, and itwas found that the iHc and squareness properties in the direction ofpressing were better than those for Example 13, and Br=12.0 kG; iHc=13.6kOe, and (BH) max=29.8 MGOe.

Example 15

The alloy ingot which was hot rolled in Example 13 was heat treated for2 hours at 1,020° C. in a vacuum, and then further heat treated for 1hour at 550*C in a vacuum. Thereafter, two pieces of this alloy ingotwere laminated together and processed to give a cube having sides of 7mm. The magnetic properties were measured by a BH curve tracer, and itwas found that the iHc and squareness properties in the direction ofpressing were better than those for Example 14, and Br=12.0 kG; iHc=14.1kOe, and (BH) max=31.6 MGOe.

Example 16

The alloy ingot of Example 6 was crushed to 500 μm or less in an argongas atmosphere using a stamp mill. The iHc measured using VSM was 2.9kOe, and degradation in the iHc was low.

Using this alloy powder, a bonded magnet having a density of 5.8 g/cm³was prepared using the same method as used in Example 7. The magneticproperties were measured by a BH curve tracer, and it was found thatBr=9.1 kG, iHc=2.8 kOe, and (BH) max=8.9 MGOe.

Example 17

The alloy ingot of Example 6 was heat treated for 5 minutes at 750° C.in a vacuum. The magnetic properties were measured by a BH curve tracer.It was found that the properties in tree directions were approximatelythe same and that the magnetic properties had improved with Br=11.8 kG,iHc=4.2 kOe, and (BH) max=15.0 MGOe.

In addition, after this magnet was magnetized, a magnetic field in theopposite direction of 2.5 kOe was applied to it, and when the magneticfield was returned to 0, it was shown that Br had a large sprig backrecovering 95% of the original, and that the magnet was an isotropicexchange spring magnet.

Example 18

The alloy ingot of Example 17 was crushed to 500 μm or less in an argongas atmosphere using a stamp mill. The iHc measured using VSM was 4.0kOe, and degradation in the iHc was low.

Using this alloy powder, a bonded magnet having a density of 5.8 g/cm³was prepared using the same method as used in Example 7. The magneticproperties were measured by a BH curve tracer, and it was found thatBr=9.1 kG, iHc=3.9 kOe, and (BH) max=92 MGOe.

Example 19

In order that the composition comprise Nd: 10.8% by atom (23.5% bymass), Pr: 3.3% by atom (7.0% by mass), Dy: 0.6% by atom (1.5% by mass),B: 6.1% by atom (1.0% by mass), Al: 0.7% by atom (0.30% by mass), Co:1.1% by atom (1.0% by mass), Cu: 0.1% by atom (0.1% by mass), Zr: 0.4%by atom (0.5% by mass), and the remainder iron, each of the startingmaterials of metal neodymium, metal praseodymium, metal dysprosium,ferro-boron, aluminum, electrolytic cobalt, electrolytic copper,ferro-zirconium, and electrolytic iron were mixed, and melted by highfrequency induction heating using an alumina crucible in an argon gasatmosphere, and then a cylindrically shaped alloy ingot was obtained bycasting under the following conditions.

The cylindrical mold was made of copper (thermal conductivity at 27° C.is 398 W/mK) and had an internal diameter of 150 mm and a length of 150mm. A film having a thickness of 100 μm and a composition of SUS304(thermal conductivity at 27° C. is 16.0 W/mK) was formed using plasmaspraying onto the surface of the inner wall of the mold. The rotatingbody was a cylindrical vessel with an internal diameter of 50 mm, and 8holes having diameters of 3 mm were formed in the side walls. Inaddition, the rotation axis L of the cylindrical mold and the rotationaxis R of the vessel were both arranged in the vertical direction.

The average rate of deposition of the molten metal on the inner wall ofthe mold was 0.05 mm/second. The number of rotations of the mold in thiscase was set to produce a centrifugal acceleration of 20 G, and the rateof rotation of the vessel form rotating body exert a centrifugal forceof approximately 10 G on the molten metal. In addition, during casting,the rotating vessel was moved back and forth over a distance of 50 mm inthe vertical direction in a 4 second cycle.

The exterior of the cylindrically shaped alloy ingot was 150 nun, thethickness at the center section along the longitudinal axis was 8 mm,and the thickness was approximately 10 mm at the thickest sections inthe vicinity of both edges. As a result of the measurement of thediameter of the crystals of the alloy ingot using a polarizingmicroscope, it was found that the area occupied by crystals having aparticle size of 10 μm or less was 96%.

A cube have sides of 7 mm was cut out from the alloy ingot and themagnetic properties were measured using a BH curve tracer The propertiesin three directions were approximately the same and the residualmagnetic flux density Br=8.1 kG, the coercive force iHc=16.8 kOe, andthe maximum energy product (BH) max=12.5 MGOe. From this, it can beunderstood that the present alloy is suitable as a cylindrically shapedisotropic magnet.

Comparative Example 8

Each of the raw materials were mixed such that the composition was thesame as the composition of Example 19, and melted by high frequencyinduction heating using an alumina crucible in an argon gas atmosphere,and casting was carried out using the same apparatus and under the sameconditions as used in Example 19. However, there was no film of any typeformed on the inner surface of the mold and the alloy was deposited andsolidified on the surface of a copper mold. The exterior of the obtainedcylindrically shaped alloy ingot was 150 mm, the thickness at the centersection along the longitudinal axis was 8 mm, and the thickness wasapproximately 10 mm at the thickest sections in the vicinity of bothedges. As a result of the measurement of the diameter of the crystals ofthe alloy ingot using a polarizing microscope, it was found that thearea occupied by crystals having a particle size of 10 μm or less was5%.

A cube have sides of 7 mm was cut out from this alloy ingot and themagnetic properties were measured using a BH curve tracer. Theproperties were highest when measured in the plane perpendicular to thesurface of the mold and were Br=2.8 kG; iHc=1.2 kOe, and (BH) max=0.4MGOe. These are extremely low compared with those of Example 4.

Example 20

Using the apparatus used in Example 19, a cylindrically shaped alloyingot manufactured using the same composition and under the sameconditions as in Example 19 was heat treated for 1 hour at 550° C. in avacuum. A cube having sides of 7 mm was cut out from this heat treatedalloy ingot. The magnetic properties were measured by a BH curve tracer,and it was found that the properties were approximately the same inthree directions with Br=8.2 kG; iHc=17.2 kOe, and (BH) max=13.1 MGOe.These magnetic properties are better than those for the alloy of Example19.

Example 21

Using the apparatus used in Example 19, cylindrical shaped alloy ingotmanufactured using the same composition under the same conditions as inExample 19 was heat treated for 2 hours at 1,020° C. in an argonatmosphere, and then further heat treated for 1 hour at 550° C. in avacuum.

A cube having sides of 7 mm was cut out from this heat treated alloyingot. The magnetic properties were measured by a BH curve tracer, andit was found that the properties were approximately the same in threedirections with Br=8.3 kG; iHc=17.5 kOe, and (BR) max=13.7 MGOe. Thesemagnetic properties are better than those for the alloy of Example 20.

Example 22

Casting was carried out using the same composition under the sameconditions as in Example 4, and using the same apparatus as used inExample 4. However, a cylindrical copper mold (thermal conductivity at27° C. is 80.3 W/mK) was used in which grooves of a depth of 1 mm with abase of a width of 5 mm had been cut at 3 mm intervals in the innersurface wall, and on which a film having a thickness of 100 μm and acomposition of SUS304 (thermal conductivity at 27° C. is 16.0 W/mK) hadbeen formed by means of plasma spraying, thereafter.

The thickness of the obtained alloy ingot was 8 mm at the center sectionof the cylindrical mold, and approximately 10 mm at the thickestsections in the vicinity of both edges. As a result of the measurementof the diameter of the crystals of the alloy ingot using a polarizingmicroscope, it was found that the area occupied by crystals having aparticle size of 10 μm or less was 98%.

A cube having sides of 7 mm was cut out from this alloy ingot. Themagnetic properties were measured by a BH curve tracer, and it was foundthat the properties were approximately the same in three directions withthe residual magnetic flux density Br=8.6 kG, the coercive forceiHc=11.0 kOe, and the maximum energy product (BH) max=14.4 MGOe.

Example 23

Using the apparatus used in Example 19, a cylindrically shaped alloyingot was obtained by mixing alloy starting materials to obtain the samecomposition as in Example 6, and using the same conditions as in Example19. However, the average rate of deposition of the molten metal on theinner wall of the mold was 0.02 mm/second.

The exterior of the obtained cylindrically shaped alloy ingot was 150mm, the thickness at the center section along the longitudinal axis was8 mm, and the thickness was approximately 10 mm at the thickest sectionsin the vicinity of both edges. As a result of the measurement of thediameter of the crystals on a cross-section of the alloy ingot using apolarizing microscope, it was found that the area occupied by crystalshaving a particle size of 1 μm or less was 65%. The area occupied bycrystals having a particle size of 1 μm or less was 65%.

A cube have sides of 7 mm was cut out from the alloy ingot and themagnetic properties were measured using a BH curve tracer. Theproperties in three directions were approximately the same and Br=11.8kG, iHc=3.0 kOe, and (BH) max=14.8 MGOe. From this, it can be understoodthat the present alloy is suitable as a cylindrically shaped isotropicexchange spring magnet.

In addition, after this magnet was magnetized, a magnetic field in theopposite direction of 2.5 kOe was applied to it, and when the magneticfield was returned to 0, it was shown that Br had a large spring backrecovering 95% of the original. It is possible to judge this magnet tobe an isotropic exchange spring magnet.

Example 24

Using the apparatus used in Example 19, a cylindrically shaped alloyingot was obtained by mixing alloy starting materials to obtain the samecomposition as in Example 6, and using the same conditions as in Example19. However, the average rate of deposition of the molten metal on theinner wall of the mold was 0.02 mm/second. In this case, the alloy ingotwas heat treated for 5 minutes at 750° C. in a vacuum. A cube havingsides of 7 mm was cut out from this heat treated alloy and the magneticproperties were measured using a BH curve tracer. The properties inthree directions were approximately the same and Br=11.8 kG, iHc=4.1kOe, and (BH) max=15.0 MGOe. The magnetic properties were better thanthose in Example 23.

In addition, after this magnet was magnetized, a magnetic field in theopposite direction of 2.5 kOe was applied to it, and when the magneticfield was returned to 0, it was shown that Br had a large spring backrecovering 95% of the original, and that this magnet was an isotropicexchange spring magnet.

Example 25

Alloy starting materials were mixed to obtain the same composition as inExample 4, and melted by high frequency induction heating using analumina crucible in an argon gas atmosphere, and then casting wasconducted under the following conditions.

The cylindrical mold was made of iron (thermal conductivity at 27° C. is80.3 W/mK) and had an internal diameter of 600 mm and a length of 600mm. A film having a thickness of 100 μm and a composition of 80% by massNi-20% by mass Cr (thermal conductivity at 27° C. is 12.6 W/mK) wasformed using plasma spraying onto the surface of the inner wall of themold The rotating body was a cylindrical vessel with an internaldiameter of 250 mm, and 8 holes having diameters of 3 mm were formed inthe side walls. In addition, during casting, the rotation axis L of thecylindrical mold was in the horizontal direction and the inclined angleθ formed by the rotation axis R of the vessel and the rotation axis L ofthe cylindrical mold was fixed at 25°. In addition, in order for removalof the alloy flakes deposited and solidified on the inner surface of themold during casting, a scraper was installed on the inner wall of themold so as to make contact with the leading edge. In addition, a metalplate guard wall having a thickness of 5 mm was arranged on the scraperside of the rotating body to prevent the direct deposition of the moltenmetal onto the scraper itself during casting.

The average rate of deposition of the molten metal onto the inner wallof the mold was 0.05 mm/second. In this example, the number of rotationsof the mold was set so that the centrifugal acceleration was 10 G, andthe rate of rotation of the vessel form rotating body exerted acentrifugal force of approximately 20 G on the molten metal.

The size of the obtained alloy flakes was of the order of 5 mm and thethickness was about 50 to 100 μm. As a result of the measurement of thediameter of the crystals of the alloy flakes using a polarizingmicroscope, it was found that the area occupied by crystals having aparticle size of 10 μm or less was 95%.

The iHc of the alloy flakes was measured by VSM and found to be 10.2kOe. This alloy flakes was crushed to 500 μm or less in an argon gasatmosphere using a stamp mill, and thereafter, a bonded magnet having adensity of 5.8 g/cm³ was prepared using the same method as used inExample 7. The magnetic properties were measured by a BH curve tracer,and it was found that Br=6.6 k; the coercive force iHc=9.8 kOe, and themaximum energy product (3H) max 8.4 MGOe.

Example 26

The alloy flakes of Example 25 was heat treated for 1 hour at 550° C. ina vacuum. The iHc of this heat treated alloy flakes was measured by VSMand found to be 10.2 kOe. This alloy flakes was crushed to 500 μm orless in an argon gas atmosphere using a stamp mill, and thereafter, abonded magnet having a density of 5.8 g/cm³ was prepared using the samemethod as used in Example 7. The magnetic properties were measured by aBH curve tracer, and it was found that Br=6.8 kG, the coercive forceiHc=10.6 kOe, and the maximum energy product (BH) max=8.9 MGOe.

Example 27

The alloy flakes of Example 25 was heat treated for 2 hours at 1,020° C.in an argon atmosphere, and thereafter it was heat treated for 1 hour at550° C. in a vacuum. The iHc of this heat treated alloy flakes wasmeasured by VSM and found to be 11.3 kOe. This alloy flakes was crushedto 500 μm or less in an argon gas atmosphere using a stamp mill, andthere after, a bonded magnet having a density of 5.8 g/cm³ was preparedusing the same method as used in Example 7. The magnetic properties weremeasured by a BH curve tracer, and it was found that Br 6.9 kG, thecoercive force iHc=11.0 kOe, and the maximum energy product (13H) max9.3 MGOe.

1. A production method for a rare earth magnet alloy ingot, wherein theproduction method comprising receiving molten metal by means of a rotarybody; sprinkling the molten metal by the effect of rotation of therotary body; and causing the sprinkled molten metal to be deposited andsolidify on an inner surface of a rotating cylindrical mold; and whereinthe inner surface includes a non smooth surface is comprised of grooveshaving a depth of 0.5 to 1 mm.
 2. A production method for a rare earthmagnet alloy ingot according to claim 1, wherein an axis of rotation ofthe rotary body and an axis of rotation of the cylindrical mold form anangle of inclination θ.
 3. A production method for a rare earth magnetalloy ingot according to claim 1, wherein the rare earth magnet alloyingot is an R-T-B magnet alloy ingot.
 4. A production method for a rareearth magnet alloy ingot comprising receiving a molten alloy of rareearth metal alloy by means of a rotary body; sprinkling the molten alloyby the effect of rotation of the rotary body, and causing the sprinkledmolten alloy to be deposited and solidify on an inner wall surface of arotating cylindrical mold; wherein a film having a thermal conductivitylower than that of material comprising the mold is provided to the innerwall surface of the cylindrical mold.
 5. A production method for a rareearth magnet alloy ingot according to claim 4, wherein the inner wallsurface of the rotating cylindrical mold is comprised of grooves havinga depth of 0.3 to 1.0 mm.
 6. A production method for a rare earth magnetalloy ingot according to claim 4, wherein the thermal conductivity ofthe film is 80 W/mK or less.
 7. A production method for a rare earthmagnet alloy ingot according to claim 4, wherein the film is made of ametal, a ceramic, or a metal-ceramic composite.
 8. A production methodfor a rare earth magnet alloy ingot according to claim 4, wherein thefilm is provided on the inner wall surface of the mold by at least oneselected from coating, plating, spray coating, and welding.
 9. Aproduction method for a rare earth magnet alloy ingot according to claim4, wherein the film has a thickness falling within a range of 1 μm to 1mm.
 10. A production method for a rare earth magnet alloy ingotaccording to claim 4, wherein an axis of rotation of the rotary body andan axis of rotation of the cylindrical mold form an angle of inclinationθ.
 11. A production method for a rare earth magnet alloy ingot accordingto claim 4, wherein two or more layers comprising rare earth alloy ingotare deposited and casted on the inner wall surface of the mold.
 12. Aproduction method for a rare earth magnet alloy ingot according to claim4, wherein the production method further comprises the step ofhot-working the obtained rare earth alloy ingot at 500 to 1,1000° C. 13.A production method for a rare earth magnet alloy ingot according toclaim 4, wherein the production method further comprises the step ofhot-treating the obtained rare earth alloy ingot at 400 to 1,000° C. 14.A production method for a rare earth magnet alloy ingot according toclaim 4, wherein the production method further comprises the steps ofheat-treating the obtained rare earth alloy ingot at 1,000 to 1,100° C.,and subsequently, heat-treating at 400 to 1,000° C.
 15. A productionmethod for a rare earth magnet alloy ingot according to claim 4, whereinthe rare earth alloy ingot is an R-T-B magnet alloy (R represents atleast one element selected from among rare earth elements, including Y;and T represents a substance predominantly comprising Fe, with a portionof Fe atoms being optionally substituted by Co, Ni, Cu, Al, Ga, Cr, andMn).
 16. A production method for a rare earth alloy flakes comprisingreceiving a molten alloy of rare earth metal alloy by means of a rotarybody; sprinkling the molten alloy by the effect of rotation of therotary body, and causing the sprinkled molten alloy to be deposited andsolidify on an inner wall surface of a rotating cylindrical mold;wherein a film having a thermal conductivity lower than that of materialcomprising the mold is provided to the inner wall surface of thecylindrical mold; and wherein casting is performed while alloy flakesdeposited on the inner wall surface of the cylindrical mold are scraped.17. A production method for a rare earth alloy flakes according to claim16, wherein the production method further comprises hot-treating theobtained rare earth alloy flakes at 400 to 1,000° C.
 18. A productionmethod for a rare earth alloy flakes according to claim 16, wherein theproduction method further comprises heat-treating the obtained rareearth alloy flakes at 1,000 to 1,100° C., and subsequently, anheat-treating at 400 to 1,000° C.
 19. A production method for a rareearth alloy flakes according to claim 16, wherein rare earth alloyflakes are R-T-B magnet alloy flakes (R represents at least one elementselected from among rare earth elements, including Y; and T represents asubstance predominantly comprising Fe, with a portion of Fe atoms beingoptionally substituted by Co, Ni, Cu, Al, Ga, Cr, and Mn).